Alloy and method for producing alloy

ABSTRACT

To provide an alloy which can suppress minute temporal deformation of a Super Invar alloy as much as possible, and a method for producing the alloy. The alloy of the present invention includes iron, nickel, and cobalt, which are the basic components of a Super Invar alloy, and is characterized in that an amount of a fraction which has not carbidized in carbon contained in the alloy is 0.010 wt % or less.

TECHNICAL FIELD

The present invention relates to an alloy with a low coefficient of thermal expansion used in a structural part of a precision device and the like, and a method for producing the alloy.

BACKGROUND ART

Conventionally, there have been publications reporting the temporal deformation of an Invar alloy (Fe and 36% Ni), which is a kind of a low thermal expansion metal material (refer to Physics and Applications of Invar

Alloys, Honda Memorial Series on Material Science No. 3, AGING AND PRECIPITATION 1978, Maruzen). This publication uses the term “γ expansion”, and conjectures that a cause for γ expansion is the carbon which is contained in the alloy.

Further, there are also examples which point out the problems of temporal dimensional change over a long period of time in the metal material used in a structural part of a precision device (refer to Japanese Patent Application Laid-Open No. H08-269613). The cause in this case is the result of a release process of the inner residual stress conferred during production processes such as a heat treatment and the like. This publication reports, although it is for a cast iron type low thermal expansion alloy with carbon content of 0.3 to 2.5 wt %, the low thermal expansion could be achieved by reducing Ni localization. Concerning the release of the inner residual stress of the latter publication, this is a known phenomenon in the production of parts made from metal, and until now procedures and treatments to suppress this phenomenon have been carried out in the production processes of precision parts.

On the other hand, regarding the former publication, this is still at the academic stage. Even among persons skilled in the art, the reality is that this phenomenon is not well known. The reason for that is considered to be that the absolute amount of the temporal deformation produced by the phenomenon of γ expansion is smaller than the temporal deformation produced by the other common phenomena.

Further, in recent precision apparatuses which are becoming substantially more precise, increasingly Super

Invar alloys, not Invar alloys, are used which reliably have a coefficient of thermal expansion at temperatures close to room temperature of less than 1 ppm (=1×10⁻⁶)/degree. Of course, while some inorganic materials have a coefficient of thermal expansion of less than 1 ppm/degree, the production processes of such materials are very difficult, such as cutting processing being impossible and the like. Further, since such materials are not very tough, the material can be damaged during the production processes. Moreover, since such materials have a small thermal conductivity, when a localized temperature distribution is produced in the member, partial expansion can occur, which prevents the characteristic of having a small coefficient of thermal expansion from being fully exploited. Therefore, there is a need to skillfully use Super Invar alloys while exploiting their characteristics.

DISCLOSURE OF THE INVENTION

However, concerning Super Invar alloys, whether a γ expansion phenomenon occurs in the same manner as in Invar alloys has not been grasped. Further, for high-precision optical apparatuses, the performance of the apparatus can gradually deteriorate over a long period of time due to a change in the optical path length as a result of temporal deformation of an important structural part. Therefore, in the investigation process leading to the present invention, to suppress the respective well-known phenomena which are causes of the temporal deformation of a metal material, Super Invar alloys were prepared in which the usual treatments were diligently performed. Despite this, temporal deformation still remained, and the amount of this deformation (temporal deformation amount of 5 ppm annually) was such that it could not be ignored as a cause for performance deterioration of optical apparatuses which will continue to become more precise in the future. If the microdeformation evaluation system which greatly contributed to the present invention is used, minute amounts of temporal deformation which have conventionally been overlooked become clear.

It is an object of the present invention to provide an alloy which can suppress as much as possible minute temporal deformation of a Super Invar alloy, and a method for producing such an alloy.

In view of the above-described problems, the alloy of the present invention includes iron, nickel, and cobalt, which are the basic components of a Super Invar alloy, and is characterized in that an amount of a fraction which has not been carbidized in carbon contained in the alloy is 0.010 wt % or less.

Further, a method for producing an alloy of the present invention is a method for producing an alloy including iron, nickel, and cobalt which are basic components of a Super Invar alloy, characterized by including:

adding carbide forming elements to the basic components and melt casting the resultant mixture;

hot forging at a predetermined temperature; and

precipitating into a base phase carbides formed from carbons contained in the alloy and the carbide forming elements by performing a first heat treatment at a first temperature which is lower than the predetermined temperature.

The present inventors discovered that minute temporal deformation of a Super Invar alloy is caused by a fraction which has not been carbidized in carbon contained in the alloy.

Further, with the alloy of the present invention, minute temporal deformation of a Super Invar alloy can be suppressed as much as possible (specifically, 2 ppm (2×10⁻⁶) or less calculated on an annual basis).

In addition, according to the method for producing an alloy of the present invention, even if only a trace amount of carbide forming elements is added, the carbide forming elements effectively combine with the carbon to form carbides. As a result, the amount of a fraction which has not been carbidized is carbon contained in the alloy can be made to be 0.010 wt % or less. Therefore, an alloy which can suppress as much as possible temporal deformation while maintaining a low coefficient of thermal expansion can be produced.

Further features of the present invention will become apparent from the following description of exemplary embodiments with reference to the attached drawings.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1A is a graph illustrating a stress-strain curve obtained as a result of a three point bending test.

FIG. 1B illustrates a schematic diagram of the three point bending test of FIG. 1A.

FIG. 2 is a graph illustrating the maximum surface stress of various composition materials and a permanent strain amount remaining at that time.

FIG. 3 illustrates an actually-measured amount of temporal deformation of the various composition materials.

FIG. 4 is a schematic diagram illustrating a measurement system used in the measurement of the amount of temporal deformation.

FIGS. 5A, 5B and 5C are photographs illustrating the metal structure of a high strength material 8-2 and a graph illustrating the results of dissolution extraction analysis.

FIG. 6 is a graph illustrating a relationship between the amount of free carbon of the various composition materials determined from the results of the dissolution extraction analysis and the temporal deformation amount.

FIG. 7 is a graph illustrating the proportion of fixed carbon content (wt %) with respect to the total carbon content (wt %), calculated using the fixed carbon content (wt %) determined from the results of the dissolution extraction analysis of FIG. 5C.

FIG. 8 is a graph illustrating a relationship between the heat treatment temperature and the amount of free carbon.

FIG. 9 is a graph illustrating the coefficient of thermal expansion of the respective Super Invar alloy composition materials used in the present investigation.

FIG. 10 is a graph illustrating the results of measuring the temperature and displacement (dimensional change of a test piece) of a high strength material (C: 0.118%) by a thermal expansion meter.

FIG. 11 is a cross sectional view of a lens barrel using a frame member 8 of the alloy of the present invention.

BEST MODE FOR CARRYING OUT THE INVENTION

The present invention will now be described in more detail with reference to the drawings. Further, the same structural elements are in principal provided with the same reference numeral, and thus a description thereof is omitted.

(Alloy)

The alloy of the present invention includes iron, nickel, and cobalt, which are the basic components of a Super Invar alloy. This alloy is characterized in that the amount of a fraction which has not been carbidized in carbon unavoidably included in the alloy is 0.010 wt % or less.

Super Invar alloys have iron, nickel, and cobalt as basic components, with a coefficient of thermal expansion which is even lower than that of Invar alloys (Fe 63.5 wt %, Ni 36.5 wt %). The basic composition of a Super Invar alloy is Fe 63.5 wt %, Ni 31.5 wt %, and Co 5.0 wt %, obtained by replacing 5 wt % of the Ni of the Invar alloy with Co.

In the present specification, “basic components” means the essential components of a Super Invar alloy.

In the alloy of the present invention, the content ranges of the above-described basic components may be illustrated as follows.

The Co content may be 2.0 wt % or more and 8.0 wt % or less, and can further be 3.0 wt % or more and 7.0 wt % or less.

The Ni content may be 30.0 wt % or more and 38.0 wt % or less. Here, when the below-described carbide forming elements, such as Ti and Nb are added in an amount higher than the minimum necessary amount to fix the carbons, excess carbide forming elements form a compound with the Ni. Thus, in such a case, a larger amount of Ni has to be included, taking into consideration the amount of Ni forming a compound with the excess carbide forming elements. Therefore, a desirable upper limit is 38.0 wt %. However, when the minimum necessary amount of the carbide forming elements is added, the upper limit is about 34.0 wt %.

The Fe content may be 62.0 wt % or more and 68.0 wt % or less, and can further be 63.0 wt % or more and 67.0 wt % or less.

The present inventors prepared a Super Invar alloy which had been subjected to common treatments for suppressing the respective well-known phenomena which are causes of temporal deformation of metal materials such as an Invar alloys and the like. Despite this, there was still temporal deformation of about 5 ppm annually. While this will be described in more detail below, the present inventors discovered that the cause of this temporal deformation is the amount of free carbon. The above-described common treatments, although will be described in more detail below, include e.g. slow cooling (furnace cooling) and a sub-zero treatment.

Here, in the present specification, non-carbide-forming non-attached carbon of carbon contained in the alloy is called “free carbon”. Specifically, a “free carbon” means a solid solution carbon such as an interstitial carbon. It could be conjectured from the results of the present experiment that free carbon is a cause of temporal deformation due to the free carbon intruding into the lattice and moving in the lattice.

If the amount of free carbon is 0.010 wt % or less, the temporal deformation could be suppressed to 2 ppm or less calculated on an annual basis. Further, the amount of free carbon is more desirably 0.005 wt % or less. A lower limit of 0.0005 wt % or less could be realized.

In the alloy of the present invention, the total carbon content in the alloy, specifically, the carbon content which is unavoidably contained in the alloy, may be 0.010 wt % or less. When a Super Invar alloy is obtained by general industrial processes, the total carbon content is often about 0.010 wt % to 0.030 wt % or less. Thus, it was learned that even if the total carbon content is 0.010 wt % or more, as long as the amount of free carbon is 0.010 wt % or less, temporal deformation can be effectively suppressed.

To ensure that as few of the carbons which are unavoidably present in the alloy as possible are not free carbon, the alloy of the present invention may include carbide forming elements. In the present specification, “carbide forming elements” are an element which is capable of combining with carbon in the alloy to form a carbide. At least a part of the carbide forming elements combines with carbon in the alloy to form a carbide, and the formed carbide is dispersed in the base phase. As long as a carbon atom is fixed as a carbide, the amount of free carbon can be decreased.

It is desirable to include a larger amount of carbide forming elements than the total carbon content based on an atomic percent comparison, otherwise free carbon which is not fixed as carbides will remain.

The carbide forming element content is desirably 0.05 wt % or more. This is because such a content allows an amount of free carbon equivalent to 0.01 wt % of the total carbon content to be fixed. On the other hand, to maintain the coefficient of thermal expansion to less than 1 ppm/° C., the carbide forming element content is desirably 0.50 wt % or less. The specific significance of these values will be described below. Further, the lower limit is more desirably 0.10 wt %. The upper limit is more desirably 0.30 wt %. The reason for this is as follows. If the carbide forming element content is too low, the carbons in the alloy cannot all be fixed as carbides, so that the temporal deformation cannot be effectively suppressed. On the other hand, if the carbide forming element content is too high, some of the carbide forming elements do not combine with the carbons and are left over, which increases the coefficient of thermal expansion. As a result, the desirable characteristics of the Super Invar alloy cannot be exploited.

The amount of carbide forming elements which do not form a carbide is desirably as low as possible. If this amount is 0.50 wt % or less, the low coefficient of thermal expansion which is characteristic of a Super Invar alloy can be maintained.

In the present invention, it is desirable that a compound phase is formed between the nickel and the carbide forming elements. Carbide forming elements which, as described above, do not form carbides and are left over, combine with the Ni, which is a basic component of the alloy. As a result, an increase in the coefficient of thermal expansion can be suppressed, and the strength of the alloy can be increased.

Examples of the carbide forming elements include titanium (Ti) and niobium (Nb).

(Lens Holding Member, Optical Apparatus)

The alloy of the present invention can suitably be used for a lens holding member. FIG. 11 is a cross sectional view of a lens barrel using a frame member 8 of the alloy of the present invention. In FIG. 11, a lens 7 made of low thermal expansion (0.6 ppm/degree) quartz is fixed on the frame member 8 having the same coefficient of thermal expansion. Further, an outer casing 9 supports the frame member.

Since the coefficient of thermal expansion is the same as the lens, light passing through the lens (the optical path) does not change even if a change in temperature occurs.

Further, since the frame member uses the alloy of the present invention, it hardly undergoes temporal deformation, so that the frame member does not impart deformation to the lens which is fixed thereto. Therefore, since optical errors, such as aberration, do not occur over a long period of time, the characteristics of the apparatus using this barrel also do not change.

Desirable examples of optical apparatuses for which the advantageous effects of the present invention would be appreciated, and which require nanometer level precision include an exposure apparatus used in manufacturing semiconductor devices, but are not limited thereto, may also used for e.g. an optical apparatus which is used in space.

(Alloy Production Method)

The alloy of the present invention can be produced as follows, for example.

If the total carbon content can be suppressed to 0.010 wt % or less, obviously the amount of free carbon will also be in this range. Although suppressing the total carbon content to 0.010 wt % or less can be achieved by exploiting sophisticated refining techniques such as VAR (vacuum arc remelting) and ESR (electroslag remelting), this is difficult with ordinary industrial process.

Further, even for a Super Invar alloy having a total carbon content of 0.010 wt % or more, by adding carbide forming elements such as Ti, Nb, and the like, melt casting the resultant mixture, and hot forging at a predetermined temperature, the carbons can be fixed by the Ti and the like, which allows an alloy with suppressed free carbon to be obtained.

However, as described above, it is desirable to avoid adding the carbide forming elements in an excessive amount, and in an amount as small as possible. Further, while it is desirable for all of the carbide forming elements to react with the carbons in just the right proportion, this is difficult in actual practice as the amount of carbons which is unavoidably present varies. Thus, it was learned that to obtain the carbides from the added carbide forming elements as efficiently as possible, the following production method may be carried out.

First, the carbide forming elements are added to the basic components, and the resultant mixture is melt cast by a 50 kg vacuum induction melting furnace.

Then, the mixture is hot forged at a predetermined temperature. The predetermined temperature may be 1000° C. or more and 1100° C. or less.

Further, by carrying out a first heat treatment at a first temperature which is lower than this predetermined temperature, the carbides formed from the carbons contained in the alloy and the carbide forming elements can be precipitated in the base phase.

The first temperature does not have to be a constant temperature, and the first heat treatment may be carried out by holding a temperature which gradually decreases. However, it is desirable to hold the temperature in a predetermined range. This is because if the temperature is too high, the formed carbide decomposes, while if the temperature is too low, it is difficult for the carbide forming elements and the carbons to bond. From such a perspective, the first temperature is desirably 825° C. or more and 950° C. or less.

The significance of this production method of the present invention is that the method does not reduce the carbon content by utilizing an expensive refining technique, but rather suppresses the undesirable effects of the carbon atoms without an increase in the coefficient of thermal expansion even if a small amount of carbon atoms is unavoidably present during a typical melt and casting production process.

Further, after the first heat treatment, the compound phase of the nickel and the carbide forming elements may be formed by carrying out a second heat treatment at a second temperature which is lower than the first temperature. As a result, excess carbide forming elements which did not contribute to carbide formation can be suppressed as much as possible from becoming a cause of an increase in the coefficient of thermal expansion, thereby allowing the strength of the alloy to be improved.

The second temperature may be 700° C. or more and 750° C. or less.

Further, after the second heat treatment, a third heat treatment may be carried out at a third temperature which is equal to or higher than room temperature to lower than the Curie temperature of the alloy. As a result, free carbon which could not turn into a carbide can diffuse to a stable position in the Super Invar alloy. By applying a so-called artificial seasoning effect, the temporal deformation can be further reduced.

The third temperature may be 25° C. or more and 150° C. or less. This upper limit may further be 120° C.

EXAMPLE (Concerning the Causes of Temporal Deformation)

First, using members of Super Invar alloys and Super Invar alloy equivalent materials, a procedure from which the causes of temporal deformation, which conventionally has been a problem, could be recursively inferred was constructed. Specifically, unlike the causes of typical temporal deformation, a hypothesis that carbon atoms have some sort of effect on temporal deformation was constructed.

The following three procedures were carried out to decrease temporal deformation: (i) decrease the amount of carbon atoms which are a cause to a level which is not a problem; (ii) prevent the unavoidably contained carbon atoms from diffusing by fixing the carbon atoms by some sort of procedure; and (iii) carry out a metal structure stabilizing treatment (so-called “seasoning”) for a practical period of time.

As a result of this investigation, it was discovered that all of these procedures were effective in reducing temporal deformation. This result is related to the fact that the carbon atoms are a cause. Further, this result indicates that containing carbon atoms in itself is not a problem, but that the problem is what state the carbon atoms are present in the metal structure. In other words, it is better not to allow the carbon atoms to diffuse in the lattice at temperatures close to room temperature.

The low thermal expansion alloy and production method thereof of the present embodiment employ what is currently thought to be the best measures. Specifically, in the melt casting step, care is taken to prevent the incorporation of impurities. The content of unavoidable carbons is suppressed, and a trace amount of an element is added for forming a carbide by strongly bonding with the carbon atoms. Further, a suitable heat treatment is carried out to effectively precipitate the carbide, so that the amount of free carbon (meaning solid solution carbons or interstitial carbons) is reduced.

The various composition materials formed from adding the carbide forming elements to the original Super Invar alloy are herein referred to as a “Super Invar alloy equivalent material”.

(Investigation of the Various Phenomena which are Causes of Temporal Deformation)

Super Invar alloys which are used for precision optical apparatus parts have a basic component elemental composition of 31.5% Ni (nickel), 5% Co (cobalt), and a balance of Fe (iron).

If the dimensional change of the material is measured at a constant temperature (23±0.01° C.) controlled with high precision, it was discovered that at an initial stage after the part was produced, extension of about 0.5 ppm calculated as a monthly basis continued. However, this is uneven depending on the production lot of the alloy, and in some lots there was greater extension. Having an assumption that the cause was due to thermally-activated process phenomena, such as diffusion and the like, the rate of change should decrease with the passage of time. Even still, it was learned from the results of measuring the dimensional change for several alloys for several tens of days, the deformation over about one month from the point where the final treatment was finished may be taken as linear change versus time elapsed, and that the extension calculated as a month could be estimated even by measurements of about 5 to 10 days.

However, it is often difficult to specify the cause of dimensional change in a metal structure based on metal structural changes. If improvements could be found by a symptomatic treatment method, whatever responded to that treatment would normally be assumed to be the cause. Even currently, where high-resolution electron microscopes are developed, micro-level changes at the atomic level can be said to be impossible to directly detect in their original state even for metal structural changes. Therefore, the cause must be conjectured based on a comprehensive consideration of information and the like from other checks, experiments, and analysis. In the present invention as well, the following respective phenomena were first investigated as the cause of this temporal deformation, and then the cause was narrowed down by a process of elimination.

(1—Concerning Reduction of Spontaneous Magnetization)

In Super Invar alloys, the same as Invar alloys, the normal atomic distance determined by the thermal vibration of the atoms is enlarged due to spontaneous magnetization specific to these alloys. As the temperature increases, the level of this spontaneous magnetization is reduced. Therefore, this spontaneous magnetization acts to cancel out the normal enlarged amount of atomic distance caused by an increase in temperature. As a result, Super Invar alloys exhibit a very small coefficient of thermal expansion at the Curie temperature or less, and at more than the Curie temperature the coefficient of thermal expansion returns to normal. Therefore, if spontaneous magnetization does decrease over time, the volume will move in a decreasing direction, which is the reverse of the presently-occurring extension (expansion) phenomenon. Thus, this phenomenon was not considered.

(2—Concerning Precipitation Phenomenon)

At close to room temperature, all of the component elements of a Super Invar alloy, Fe, Ni, and Co have a very small self-diffusion coefficient, and thus the diffusion of these atoms would be difficult to consider.

However, while atoms such as carbon which have a small atomic radius do exhibit interstitial diffusion even at room temperature, from the results of diffusion, it cannot be confirmed whether some kind of stable phase (e.g., graphite, cementite and the like) is generated.

Further, because this is an important point it will be described in more detail below, but concerning temporal dimensional change, in the range of from room temperature, which is the Curie temperature or less, to 210° C., reversible expansion, which was not normal thermal expansion, and contraction were found. It cannot be denied that there is a possibility of precipitation even at such a low temperature range. However, that precipitated material is unconceivable to re-dissolve and again form a solid solution. Therefore, since this phenomenon is not reversible at the above-described temperature range, it was excluded as the cause of temporal deformation.

(3—Concerning Release of Inner Residual Stress)

This phenomenon is a process which has inner stress as a driving force, in which as a result of the movement and disappearance of dislocations, which are a kind of lattice defect, dislocation density decreases. Thus, volume can be thought to move in a decreasing direction. However, just in case, a temporal deformation comparison measurement was carried out, using the same materials, between a material obtained by air cooling from 315° C. and a material obtained by furnace cooling, yet no difference was found. Therefore, this phenomenon was also excluded.

(4—Concerning Martensitic Transformation)

This phenomenon involves an increase in volume. However, the unique, high-precision dimensional change measurement performed in the present invention was carried out in a location which was controlled at ±0.01° C. of 23° C. However, for the present alloy, it can be considered that this transformation will not occur at a constant temperature. This is because, unlike typical tool steel and the like, the stable phase at room temperature for the present nickel-containing, iron-based low thermal coefficient alloy is originally an austenite phase. Further, there is no change in stress applied on the alloy during the dimensional change measurement. Therefore, stress-induced transformation also does not progress. However, despite this, using the same materials, a material which was subjected to a treatment at a low temperature (sub-zero treatment) of −10° C. and a material which was not subjected to such a treatment were compared. From the results, no difference was found in the temporal deformation amount. Therefore, this phenomenon was also excluded.

(5—Concerning Creep)

Considering the amount of temporal deformation which is a problem, creep deformation caused by the sample's own weight may also contribute. For typical steel, which generally is a body-centered cubic lattice, the diffusion of interstitial elements tends to occur even at room temperature, so that strain aging dynamically progresses, which makes it difficult for creep deformation to occur. Concerning this point, for this alloy which is a face-centered cubic lattice, the diffusion rate of carbon is smaller than that for steel, so that it is said that strain aging at temperatures close to room temperature does not progress easily, and creep occurs more easily.

Accordingly, the strength in the microstrain range (at a permanent strain of 2 ppm=about 0.0002%) of the various alloys was measured. It can be said that dislocations are less likely to diffuse and creep resistance is larger as the strength of the alloy at this strain amount is larger. Among the samples which were measured for temporal deformation, as expected the high-carbon material (C: 0.118 wt %) had a larger strength than the low-carbon material (C: 0.002 wt %). Thus, if creep is the cause, high-carbon materials should have a smaller amount of temporal deformation. Here, in the measured samples, other than carbon, the contained components were exactly the same. Various samples were also prepared by melt production from both electrolytic iron and electrolytic nickel, and then further adding carbon for the high-carbon materials. The results of the temporal deformation measurement will be described below.

(6—Concerning γ Expansion)

The gaps in an austenite phase (γ phase) face-centered cubic (fcc) lattice are the centers of an octahedron and the centers of a tetrahedron, the latter one having geometrically even smaller gaps. For example, if carbon atoms which were in gaps in the center position of an octahedron diffuse into gaps in the center position of a tetrahedron the volume may expand. Due to the existence of spontaneous magnetization, carbon atoms can stably be present even in positions where the geometric gap is small. If the mechanism of γ expansion is based on diffusion between nearby gaps among the lattices of carbon atoms such as described above, then the temporal deformation amount can be expected to decrease for a low-carbon material having few free carbon. Further, since the deformation is not produced by production of a compound or the like, the mechanism can be expected to be reversible with respect to the temperature change of the temperature range mentioned above. While this will be described in more detail below, based on a series of investigations for the present invention, the cause of temporal deformation could be specified as being this phenomenon.

(Alloys According to the Examples)

Table 1 is a list of various Super Invar alloys and Super Invar alloy equivalent materials which were prepared for measuring the temporal deformation amount.

TABLE 1 Prediction of Separate Various Component (wt %) Causes of Temporal Composition Total Other Main Deformation Materials Carbon Compo- Creep γ Expansion Name Code Content nent (s) Hardness Ease Ease Comparative Conventional STD 0.012% About 120 HV This This Example 1 Material conventional conventional material is material is the standard the standard Example 1 Low-Carbon LC 0.002% About 120 HV Easy for Difficult Material creep to for γ occur expansion to occur Comparative High-Carbon HC 0.118% About 120 HV Difficult Easy for γ Example 2 Material for creep to expansion to occur occur Example 2 High 8-1 0.013% Ti 3.0% Equivalent Very Difficult Strength DA to 372 HV difficult for γ Material for creep to expansion to occur occur Example 3 High 8-2 0.013% Ti 3.0% Equivalent Very Difficult Strength STA to 389 HV difficult for γ Material for creep to expansion to occur occur Example 4 High 9-1 0.012% Nb 3.9%, Ti Equivalent Very Difficult Strength DA 0.6% to 287 HV difficult for γ Material for creep to expansion to occur occur Example 5 High 9-2 0.012% Nb 3.9%, Ti Equivalent Very Difficult Strength STA 0.6% to 252 HV difficult for γ Material for creep to expansion to occur occur Example 6 Nb-Added N 0.017% Nb 0.24% Difficult Difficult Material for creep to for γ occur expansion to occur Example 7 Ti-Trace T 0.023% Ti 0.08% Difficult Difficult Amount Added for creep to for γ Material occur expansion to occur

Table 1 simultaneously illustrates, in a case that the cause of temporal deformation was creep, whether it is predicted that the various materials are not as susceptible to temporal deformation as the conventional material, or it is predicted that the various materials are more susceptible to temporal deformation than the conventional material. Table 1 also similarly illustrates these cases for when γ expansion was the cause.

Although the contained carbon content for Super Invar alloy equivalent materials thoroughly containing Ti and Nb (high strength materials, Examples 2 to 5) was basically the same as for the conventional material, the carbons would be expected to be carbides such as TiC and NbC would be expected (the analysis results will be described below). Thus, since the number of carbon atoms which can actually move should be smaller, the temporal deformation would be expected to decrease in either case that the cause is creep or γ expansion.

Hardness is written as an average value, because a hardness of “about 120 HV” depends on the location of the hardness test. Further, “equivalent to 372 HV” is actually a value obtained by converting the value of the results found by testing with a Rockwell hardness testing machine to a Vickers hardness.

In addition, the DA material of the various composition materials means a material obtained by hot forging, then air cooling, and then performing an aging treatment. Further, an STA material means a material obtained by hot forging, then performing a solution treatment, then water cooling, and then performing an aging treatment.

(Method for Producing the Alloys According to the Example)

The conventional material, low-carbon material, high-carbon material, and Nb-added material (Examples 1, 6, Comparative Examples 1 and 2) were produced by undergoing melt casting by a vacuum induction melting furnace, air cooling after hot forging at 1000° C., water cooling after a solution treatment by holding at 830° C. for 2 hours, performing a stress release treatment by holding at 315° C. for 3 hours, and performing a stabilizing treatment at 98° C. for 48 hours (so-called artificial seasoning, the third heat treatment).

Further, the high-strength materials (Examples 2 to 5) were produced by undergoing melt casting by a 50 kg vacuum induction melting furnace, hot forging at 1000° C., performing an aging treatment at 720° C. for 6 hours (second heat treatment), then performing a stress release treatment at 315°, and performing a stabilizing treatment at 98° C. for 48 hours.

Further, the Ti-trace added material (Example 7) was produced by undergoing hot forging at 1000° C., holding at 900° C. for 2 hours (first heat treatment), then gradually cooling to 830° C., air cooling from that temperature to room temperature, then performing an aging treatment at 720° C. for 6 hours (second heat treatment), then performing a stress release treatment at 315°, and performing a stabilizing treatment at 98° C. for 48 hours (third heat treatment).

(Contained Components of the Alloys According to the Example)

Table 2 lists the contained components of the various composition materials.

TABLE 2 Provided Test Materials in the Present Invention Weight (%) C Si Mn P S Cu Ni Cr Comparative Conventional 0.012 0.24 0.33 0.003 0.004 <0.01 32.69 — Example 1 Material Example 1 Low-Carbon 0.002 0.25 0.34 <0.002 0.001 <0.01 32.8 <0.01 Material Comparative High-Carbon 0.118 0.25 0.35 <0.002 0.001 <0.01 32.8 <0.01 Example 2 Material Example 2 High Strength 0.013 0.26 0.36 <0.01 <0.01 <0.01 36.9 0.01 DA Material 8-1 Example 3 High Strength Material 8-2 Example 4 High Strength 0.012 0.25 0.35 <0.01 <0.01 <0.01 36.9 <0.01 DA Material 9-1 Example 5 High Strength Material 9-2 Example 6 Nb-Added 0.017 0.25 0.25 0.005 0.012 <0.01 33.76 0.02 Material Example 7 Ti-Trace 0.023 0.25 0.35 <0.002 0.001 <0.01 32.59 <0.01 Amount Added Material Weight (%) Mo Co Ti Nb O N Fe Comparative Conventional — 4.86 — — Balance Example 1 Material Example 1 Low-Carbon <0.01 4.84 — — Balance Material Comparative High-Carbon <0.01 4.92 — — Balance Example 2 Material Example 2 High Strength — 4.88 3.0 — Balance DA Material 8-1 Example 3 High Strength Material 8-2 Example 4 High Strength — 4.93 0.6 3.9 Balance DA Material 9-1 Example 5 High Strength Material 9-2 Example 6 Nb-Added — 3.75 <0.01 0.24 Balance Material Example 7 Ti-Trace — 5.10 0.08 — 0.0008 0.0006 Balance Amount Added Material

(Conventional Material, Low-Carbon Materials, High-Carbon Materials)

The low-carbon materials and high-carbon materials were each produced using high-purity materials produced by electrolysis as the raw materials by changing only their carbon contents, and having identical amounts for the other elements.

(Concerning the High-Strength Materials)

Each of the high-strength materials had the same carbon content as the conventional material, and was made to thoroughly contain Ti and/or Nb.

The results of the hardness testing were that after the hot forging, materials 8-1 and 9-1 respectively had a hardness of 36.6 HRC (equivalent to 360 HV) and 25.7 HRC (equivalent to 270 HV), and that after the solution treatment, materials 8-2 and 9-2 respectively had a hardness of 33.3 HRC (equivalent to 330 HV) and 18.9 HRC (equivalent to 234 HV). Since the hardness illustrated in Table 1 is a value after the aging treatment (720° C.×6 hours), all of the hardness values were increased by carrying out the aging treatment for precipitation phase formation.

Further, Ti addition had a larger effect on increasing hardness than Nb addition. Ni forms a precipitation phase, which is a compound phase with the Ti and Nb, by the aging treatment. Therefore, the Ni content is increased based on the carbide forming element content so that the matrix portion becomes a composition which can exhibit an Invar effect (effect in which the coefficient of thermal expansion decreases) to the maximum extent.

(Concerning the Nb-added Material)

An Nb-added material is a material obtained by adding a small amount (0.24%) of Nb to a conventional material. The purpose of this material is to reduce γ expansion, which is said to be a cause of the above-described diffusion of carbon, by making unavoidably-contained carbons bond with the Nb, whereby the carbons are fixed as a carbide.

The atomic weights of niobium and carbon is respectively 92.9 and 12. As described below, the niobium carbide confirmed in the present invention is NbC, and thus the atoms bond at a 1:1 ratio. Therefore, unless 7.74 times or more (=92.9/12) of niobium than carbon is added, free carbon will remain. Further, since it is difficult for all of the Nb to bond with the carbon atoms, it is desirable to add more than 7.74 times of niobium than carbon. However, the addition of more Nb than necessary amount to form the carbides should be suppressed as much as possible. This is because residual Nb which does not form a carbide remains, and if this remaining niobium forms a solid solution in the matrix, the coefficient of thermal expansion is increased. The same can be said for the carbide forming elements other than Nb.

In the present example, since the carbon content is 0.017 wt %, it is desirable to add 0.13 wt % or more of Nb. Thus, the Nb content was adjusted to 0.24 wt %.

Even in this case, like with the above-described high strength materials, it is desirable to carry out an aging heat treatment (second heat treatment) on the niobium formed in a solid solution. This is because the hardness is increased due to the formation of a precipitation phase of Nb and Ni. This aging heat treatment is carried out to increase the strength value at the microstrain (micro yield point). This is because the present invention is directed to resolving the problem of a temporal deformation which is very small. The strength value at the microstrain is the stress of a remaining permanent strain at a level of 0.0001% (1 ppm). If this micro yield point can be increased, creep resistance also increases. At the above-described production conditions, this aging heat treatment is equivalent to 720° C. While the measurements here measured the temporal deformation in a state where a low thermal expansion alloy was arranged on a flat surface, in an actual apparatus, a larger stress than this state may be applied on a member which is formed from a low thermal expansion alloy. Therefore, it is even more desirable to improve creep resistance. Further, increasing the hardness of the alloy by even a little bit achieves desirable results also in the cutting and grinding processing. This is because it is easier to remove the swarf and clogging of the grinding stone is reduced.

Obviously, not only is the γ expansion of the respective high strength materials reduced, but an effect in suppressing creep deformation can also be expected. This is because the strength value at the microstrain is increasing as a result of the precipitation hardening and the solid solution hardening by the carbide forming element.

(Concerning Ti-Trace Amount Added Material)

This material has the minimum necessary amount of Ti added thereto in consideration of a case where ideally all of the carbons unavoidably present in the alloy are used in the formation of titanium carbides.

Since Ti in solid solution reduces spontaneous magnetization, and effectively increases the coefficient of thermal expansion, the minimum necessary amount of Ti to maintain a low coefficient of thermal expansion was used.

However, while as a result the temporal deformation could be reduced by a considerable amount, the component blend could not be adjusted to a Ti value which was just enough. Namely, since the carbon content was 0.023 wt %, even if all of the added 0.08 wt % of Ti ideally formed into carbides, the Ti was still not quite enough. Specifically, since the atomic weight of titanium is 47.9, at least 0.003 wt % (=0.23−0.08×12/47.9) of free carbon, which admittedly is a tiny amount, will nonetheless remain.

By carrying out a suitable carbide forming heat treatment (first heat treatment; in the experiment results, 825° C. or more and 950° C.), a considerably large proportion of 74% of the contained Ti becomes a carbide at the optimum temperature of 900° C., so that 0.008 wt % of free carbon remained. With about this amount of free carbon, the actual temporal deformation could be reduced as described below.

Further, the coefficient of thermal expansion also shows almost no difference depending on the cutting direction of the test piece. This value was 0.42 to 0.47 ppm/degree in the range of 18 to 28° C. This measurement was carried out using the TMA 8310 manufactured by Ulvac-Riko, Inc., at a rate of temperature increase of 5° C./minute.

Further, with the normally used dissolving and casting process, about 0.03 wt % of carbon is unavoidably contained. To reduce this level to 0.008 wt % is difficult in actual practice. By adding a trace amount of Ti, the amount of carbon which influences the temporal deformation could be reduced to 0.008 wt %, while hardly changing the coefficient of thermal expansion and with the same level of production costs.

(Specifying the Main Causes of Temporal Deformation)

FIG. 1A is a graph illustrating a stress-strain curve obtained as a result of a three point bending test. In FIG. 1A, the vertical axis represents the center concentrated load (kN), and the horizontal axis represents the maximum strain (ppm). Further, in FIG. 1A, the circles represent a measurement point when adding the load, and the triangles represent a measurement point when removing the load. FIG. 1B illustrates a schematic diagram of the three point bending test. Since it is the microstrain portion which is of consequence in creep deformation, the portions with a large strain are not illustrated. In FIG. 1A, as an example of the conventional material, an increase in the permanent strain amount remaining while the maximum center load is gradually increased is illustrated.

As the test piece, the various composition materials illustrated in Table 1 were prepared in a 30×30×339 mm size. The test piece was supported near either end with a span interval of 250 mm. A concentrated load was applied on a point in the center section, and the strain amount was determined from the output of a strain gauge stuck to the center section of the opposite surface. The testing temperature was room temperature, and the subsequent temperature when measuring the temporal deformation was roughly the same. The loading velocity was set to a small level of 0.5 mm/minute so that the effects of elastic after-effects could be ignored.

First, the maximum load was set to a small level, and the load was repeatedly added and removed. Then, while gradually increasing the maximum load, the maximum load at which remaining permanent strain could initially be detected was 12 kN. The remaining permanent strain (a) at that stage was 3 ppm. However, in the present test, since the strain gauge is attached to a part of the test piece where the maximum strain is produced, and the strain amount is measured at that portion, the average strain amount of the whole test piece will be an even smaller value.

Subsequently, the remaining permanent strain (b) when the maximum load was set to 15 kN, that load was applied, and then the load was removed, was 10 ppm. In this case, if the maximum load of 15 kN was applied at the start, a permanent strain of 13 ppm should remain.

FIG. 2 is a graph illustrating the maximum surface stress of the various composition materials of Table 1 and the permanent strain amount remaining at that stage at that location. In FIG. 2, the maximum load was converted into the maximum surface stress from the maximum load determined by the three point bending test of FIG. 1A and the permanent strain amount remaining at that stage, and the permanent strain amount used the cumulative value of the remaining strain amount produced until that stress.

Typically, the permanent strain amount used when displaying a strength value is 0.2% (200 ppm). However, in the present investigation, the strength value for a very small permanent strain amount was determined. In high-precision apparatuses, the deformation amount which influences the performance of the apparatus is, considering the use period and maintenance period as the lifecycle of the apparatus, about 10 ppm for the initial year. Thus, the measurement was carried out having a resolution of 1 ppm, which is one order smaller.

From these results, if the cause of temporal deformation is the phenomenon of creep deformation, the high strength materials (Examples 2 to 5) should have a markedly smaller temporal deformation. Further, for the three test pieces on the left of FIG. 2, in which only the carbon content is different (low-carbon material: Example 1; conventional material: Comparative Example 1; and high-carbon material: Comparative Example 2), the strength value increased as the carbon content increased. As a result, similarly, supposing that temporal deformation is produced by creep deformation, the amount of temporal deformation should decrease for high-carbon strength materials.

FIG. 3 illustrates the actually-measured amount of temporal deformation of the various composition materials. In FIG. 3, the amount of temporal deformation per month is illustrated for when the various nickel-containing, iron-based low thermal expansion alloys of Table 1 are assumed to undergo linear deformation for 1 month. From the results, stating the conventional material(Comparative Example 1) as a standard, the high-carbon material (Comparative Example 2) showed very large amounts of temporal deformation. Thus, a comparison among materials which have different carbon contents showed that the amount of temporal deformation increased for the high-carbon materials. As long as the high-carbon material was produced in an ordinary manner, the alloy composition contained a large amount of carbon which would be impossible as the unavoidable carbon content. To further pursue the causes of temporal deformation, along with a low-carbon material, alloys were specially produced which had a very low and a very high carbon content. In a comparison among the materials with different carbon contents (Example 1 and Comparative Examples 1 and 2), the amount of temporal deformation increased for the high-carbon materials.

Further, as will be described below, it was discovered that size hysteresis clearly appeared in the high-carbon materials during the heating and cooling when measuring the thermal expansion.

On the other hand, the respective high strength materials (Examples 2 to 5) all had a very small temporal deformation. The high strength materials contained an amount of about 3.0 wt %, which was considerably larger than the minimum necessary amount to form the carbides, of carbide forming elements. Therefore, the strength value should be increased, there should be a marked effect on reducing the temporal deformation for which creep is a cause, and there should also be a marked effect on reducing the temporal deformation for which γ expansion relating to the carbon atoms is a cause.

Therefore, while it could not be specified whether the cause was creep deformation or γ expansion in a comparison of the results of the temporal deformation for the conventional material and the high strength materials, considering the results of the above-described comparison among the materials in which only the carbon content was changed, the main cause of temporal deformation can be presumed to be γ expansion.

The Nb-added materials and the Ti-trace amount added materials were also considerably improved compared with the conventional material. Even still, since the amount of free carbon is slightly more than the high strength materials, the temporal deformation was slightly larger than that for the high strength materials. For the conventional material (Comparative Example 1) and Examples 2 to 7, even though the total carbon content was about the same, the amount of free carbon for the latter was reduced. Further, the temporal deformation in Examples 2 to 7 was quite suppressed. From these results, it was learned that the amount of the temporal deformation has a close relationship with the amount of free carbon, and not with the total carbon content.

Among the common materials used for comparison, quartz produces hardly any temporal deformation. Further, stainless steel (SUS 316, 650° C.×2 hours, then furnace cooling) and carbon steel (S45C, 800° C.×2 hours, then furnace cooling) had a negative value, and thus it can be considered that a slight inner residual stress remained. It is considered that γ expansion does not occur in either of these materials.

FIG. 4 is a schematic diagram illustrating a measurement system used in the measurement of the amount of temporal deformation. In this measurement system, a test piece 1 has the same dimensions as previously used in the three point bending test. Mirrors 4 and 5 are attached to either end of the test piece 1. These mirrors 4 and 5 are provided to reflect the laser light from two laser interferometers (end-measuring machines) 2 and 3. The laser interferometers 2 and 3 read a dimensional change in the test piece based on the change in the optical path length, and use the principles of a so-called “Michelson interferometer”.

Expansion and contraction of the test piece is determined by the difference in the positional change of either end of the test piece as measured by the two laser interferometers. The resolution of this measurement system is 0.2 nm. Since the length of the test pieces is 339 mm, conversion shows that a resolution of 0.0006 ppm can be obtained. With such a resolution, the precision is such that measurement can be sufficiently performed even with a change of about 1 ppm per year.

This measurement system is built on a quartz platen 6. The measurement system is set in an indoor environment with a temperature controlled within a range of ±0.01° C. from room temperature at 23° C. In actual practice, the temperature change of the test pieces is even smaller. As a result, the amount of expansion and contraction due to thermal expansion can be ignored.

(Method for Measuring the Amount of Free Carbon)

FIGS. 5A to 5C are photographs illustrating the metal structure of the high strength material 8-2 as an example and a graph illustrating the results of dissolution extraction analysis. By subjecting the surface to a smoothing treatment into a mirror surface, and then performing a suitable etching treatment to match the objective, each of the phases can be distinguished. In the metal structure, other phases different from the base phase are clearly dispersed. FIG. 5A is a photograph of one of those phases observed by scanning electron microscopy. A phase of about 5 μm in size can be confirmed in the center of the photograph. This phase was subjected to analysis by an Electron Probe Micro Analyser (EPMA), and titanium and carbon were detected.

Next, the method for carrying out the dissolution extraction analysis performed here will be described. First, using 0.5 g of the alloy material as the anode and platinum as the counter electrode, electrolysis was carried out at a voltage of 0.2 V in an electrolysis solution which mixed 10 v/v % of acetylacetone and 1 w/v % of tetramethylammonium chloride. Then, the resultant residue was subjected to suction filtration in which the residue was passed through polycarbonate type (PC) filter paper (0.2 μm filter) to capture the residue. FIG. 5B is a photograph obtained by observation with scanning electron microscopy in the same manner of the product captured on this filet paper. When this product was analyzed by EPMA in the same manner as described above, it was learned from the results of FIG. 5C that the product was TiC.

Next, the filter paper was ashed along with the captured product in a platinum crucible. The resultant product was then charged with a mixture (solvent) of sodium borate and sodium carbonate, and the resultant mixture was made to dissolve at 900° C. Then, the mixture was charged with 10 mL of hydrochloric acid and a few mL of perchloric acid. The dissolved solution was diluted to a constant volume, and then subjected to IPC analysis. The proportion of carbides in the dissolved alloy with respect to the alloy total amount was thus determined, and the mass of carbon was calculated from the respective atomic weights of Ti and C. Further, the amount of free carbon was determined by subtracting the mass of carbon formed as carbides from a total carbon content analyzed beforehand.

Here, since the nitrogen in the alloy may also form TiN, a separate analysis was carried out. However, in this alloy TiN was not detected. Further, while the high strength materials (9-1 and 9-2, Examples 4 and 5) simultaneously contained Ti and Nb, the carbon content formed as carbides was determined from the respective contents and atomic weights. This can be determined in the same manner even if the carbide forming elements are Ta, Zr and the like.

The total carbon content was determined by a combustion infrared-absorbing method. This analysis is a method in which a sample is heated to a high temperature in an oxygen stream to oxidize the contained carbon into carbon dioxide and the like, and the carbon content is determined by measuring the infrared absorption.

(Relationship Between the Amount of Free Carbon and Temporal Deformation)

FIG. 6 is a graph illustrating the relationship between the amount of free carbon of the various composition materials determined from the results of the above-described dissolution extraction analysis and the temporal deformation amount. In the graph, the composition materials with the temporal deformation concentrated near the origin were five composition materials, the low-carbon material (Example 1) and various high strength materials (Examples 2 to 5). The other plots also similarly correspond to Table 1 and FIG. 3.

At the plots with a low amount of free carbon, there was a direct relationship between temporal deformation amount and the amount of free carbon.

Further, the various high strength materials had a large portion of their contained carbon fixed as carbides. The proportion of the amount of free carbon to the total carbon content was 0.0031 wt % or less for all of those materials. Some of the materials had a proportion which was smaller than that of the low-carbon material (Example 1, C: 0.002 wt %), which were produced from special raw materials, such as electrolytic iron, electrolytic nickel or the like, to have a very low carbon content. As illustrated in FIG. 7, the amount of free carbon in the high strength materials 8-1 (Example 2), 8-2 (Example 3), 9-1 (Example 4), and 9-2 (Example 5) was respectively 0.0008 wt % (=0.013−0.0122, hereinafter the same), 0.0005 wt % (=0.013−0.0125), 0.0031 wt % (=0.012−0.0089), and 0.0025 wt % (=0.012−0.0095). Further, similarly, the amount of free carbon in the Ti-trace added material (Example 6) and Nb-added material (Example 7) was respectively 0.0080 wt % (=0.023−0.015) and 0.0069 wt % (=0.017−0.0101).

These materials also had a very small amount of temporal deformation. It was learned that if the amount of free carbon is made to be 0.010 wt % or less, temporal deformation can be suppressed to a level lower than conventional levels. More desirably, to suppress the amount of temporal deformation to 1 ppm or less calculated on an annual basis, the amount of free carbon is desirably 0.0050 wt % or less.

Since in the conventional material the carbon is present as free carbon, unevenness occurs in the amount of temporal deformation due to unevenness in the unavoidable carbon content which is included during production. Even for a relatively good production lot, an expansion of 0.4 ppm calculated on a monthly basis was shown. In the present invention, temporal deformation can reliably be reduced to less than that of the conventional material by causing carbides to form.

FIG. 7 is a graph illustrating the proportion of fixed carbon content (wt %) with respect to the total carbon content (wt %), calculated using the fixed carbon content (wt %) determined from the results of the dissolution extraction analysis of FIG. 5C. From the graph, Ti (Examples 7, 2, 3) tends to bond with the carbon atoms more effectively than Nb (Examples 6, 4, 5), so that only a small amount of the carbide forming elements needs to be used.

Since excess carbide forming elements causes the coefficient of thermal expansion to increase, from this perspective Nb is better than Ti. Further, in a comparison among composition materials similarly added with Nb, Examples 4 and 5, which were subjected an intermediate heat treatment (second heat treatment) at 720° C. for 6 hours, formed the carbides at a slightly larger proportion than Example 6, which was not subjected to an intermediate heat treatment. Therefore, to make the amount of temporal deformation small and to suppress the coefficient of thermal expansion to a low level, it is desirable to add the minimum necessary amount of carbide forming elements, and to carry out a suitable heat treatment in the carbide formation.

As described above, in the present experiment, in the

Super Invar alloys, it was learned that the carbide forming element having the highest carbide forming performance is Ti. However, other elements commonly mentioned as carbide forming elements in steel materials, such as Ta, Zr, W, V, Mo and the like, can also be considered to provide the effects of the present experiment.

Further, while the intermediate heat treatment conditions in the present experiment were 720° C.×6 hours, these intermediate heat treatment conditions were carried out on only the four high strength materials along with the aging treatment conditions. The primary objective of this intermediate heat treatment was to harden the materials by causing a compound phase of the Ni and the carbide forming elements to precipitate.

(Desirable Method for Producing the Alloy)

To reliably reduce the amount of free carbon, like the four high strength materials, a large amount of the carbide forming element may be added. However, if this carbide forming element is left over, the coefficient of thermal expansion is increased. Therefore, it is desirable to reduce the amount of free carbon by effectively combining as small an amount as possible of the carbide forming elements with the free carbon. To achieve this, it is desirable to carry out melt casting, hot forging, and then perform a predetermined heat treatment (first heat treatment).

FIG. 8 is a graph illustrating a relationship between the heat treatment temperature and the amount of free carbon. This graph was prepared in order to determine the temperature where the carbides tend to form in the first heat treatment. Using the Ti-trace amount added alloy (Example 7), the amount of free carbon in a total of six samples which had been melt cast by a vacuum melting furnace, hot forged at 1000° C., and then subjected to a heat treatment by holding at a predetermined temperature (first temperature) for 2 hours, and one sample which had been subjected to a solution treatment at 1100° C., was measured. As the first temperature, the measurement was carried out at intervals of 25° C. in a temperature range of 825° C. to 950° C.

From the results of the present experiment, it was learned that the first temperature at which the carbides tend to form the most and the amount of free carbon is reduced is 900° C. For example, a sample which was heat treated at 900° C. had a total carbon content of 0.023 wt %, of which the fixed carbon content was 0.0148 wt %. Specifically, 74% of the total carbon content was formed as carbides.

Further, since the Ti-trace added material had a Ti content of 0.08 wt % and carbon content of 0.023 wt %, even if all of the Ti ideally combined with the carbons, consideration should be given to the fact that 0.0030 wt % of free carbon will still remain.

From the present experiment, it was learned that the carbides form more easily by carrying out the first heat treatment at a temperature between at least 825° C. or more and 950° C. or less, and more desirably at 875° C. or more and 925° C. or less.

(Various Measurement Results)

FIG. 9 is a graph illustrating the coefficient of thermal expansion of the respective Super Invar alloy composition materials used in the present investigation. The coefficient of thermal expansion of the Super Invar alloy composition materials is the data obtained simultaneously with the results of the dimensional change measurement carried out by the thermal expansion measurement device (Laser Thermal Expansion Meter LIX, manufactured by Ulvac-Riko, Inc.) used to examine the relationship between temperature and displacement. The test piece has a size of 6 mm in diameter and 12 mm in length. The heating and cooling rates were set to 1° C./minute. Since a slight difference occurs between the temperature of the actual test piece and the temperature of a thermocouple arranged near to the test piece in order to measure the atmosphere temperature, the heating and cooling was carried out in such a slow manner. The measurement range was between 30° C. and 50° C.

As a result, as the amount of solid solution carbon increased, the coefficient of thermal expansion increased, and as the content of the carbide forming elements increased, the coefficient of thermal expansion increased.

Further, no difference was found in the coefficient of thermal expansion between the DA materials and STA materials. The STA materials were produced to have a uniform metal structure by rapidly cooling after carrying out a solution treatment. These STA materials were produced in order to minimize the Ni segregation in the metal structure as a Super Invar alloy. This is because if portions are present in the metal structure which have a low Ni content or a high Ni content, those portions will be excluded from the composition which was originally meant to have the lowest coefficient of thermal expansion, so that as a result the coefficient of thermal expansion of the whole material is increased. The influence that the respective elements have on the coefficient of thermal expansion depends on what state those elements are present in. While this is the same even for other carbide forming elements such as Nb, for example, even if Ti is added, it is desirable for the added Ti to bond with the carbon atoms which are unavoidably present, and thus be present in the metal structure as TiC. This is because, strength increases, and since the TiC itself has a small coefficient of thermal expansion, coherence with the base phase is small, which means that the action to increase the coefficient of thermal expansion of the whole material is small.

However, if the Ti atoms which do not bond with the carbon atoms are present by replacing the original lattice points of the Super Invar alloy, the Invar effects are reduced. In any event, the strength of spontaneous magnetization decreases due to the carbons and the carbide forming elements being formed in solid solution. Ti atoms and the like which do not contribute to the bonding with the carbon atoms become surplus atoms for a Super Invar alloy which aims to have a low thermal expansion.

Accordingly, to achieve a low coefficient of thermal expansion, the surplus elements may be subjected to a suitable aging heat treatment to increase strength, so that precipitation phases such as a γ′ (gamma prime) phase and a γ″ (gamma double prime) phase are generated. This is because simultaneously creep resistance can be improved and cutting processability can also be improved due to stickiness being decreased.

In a Super Invar alloy, unless a special element is added and a suitable heat treatment is carried out, the carbons which are unavoidably present will be present as free carbon (solid solution carbons or interstitial carbons). Those carbon atoms promote temporal deformation and also increase the coefficient of thermal expansion.

Concerning how much the carbide forming elements increase the coefficient of thermal expansion, in the present investigation, the increase when 0.24 wt % of Nb was added was 0.25 ppm/° C., and the increase when 3.9 wt % of Nb was added was 3.0 ppm/° C. For a 36% Ni Invar, the coefficient of thermal expansion is usually 1 ppm/degree. Since this is too high, Super Invars were originally selected. To keep a low coefficient of thermal expansion (less than 1 ppm/degree), which is the greatest characteristic of Super Invars, the above-described value of 0.50 wt % or less was determined from the approximation curve obtained during the processes of the present investigation which is illustrated in expression (2), as a condition for having a coefficient of thermal expansion of less than 1 ppm/degree.

y=0.76x+0.63  (2)

Here, y denotes the coefficient of thermal expansion in units of ppm/degree, and x denotes the Nb content in units of wt %.

For Ti, the coefficient of thermal expansion can be determined from the approximation curve illustrated in expression (3). Here, x is the same value as for Nb.

y=0.71x+0.60  (3)

While Ti has a slightly larger ratio which increases the coefficient of thermal expansion, the value is roughly the same.

FIG. 10 is a graph illustrating the results of measuring the temperature and displacement with the above-described thermal expansion meter when the high-carbon material (Comparative Example 2, C: 0.118%) was heated and cooled at a rate of 1° C./minute.

In the measurement results of the first cycle, a characteristic was that the slope of the thermal expansion curve of the test piece decreased at around 150° C. This is thought to be due to the diffusion of the carbon atoms becoming active at this temperature. If the temperature is heated to 210° C. (205° C. in the third and fourth cycles) and then cooled to room temperature, clearly different trajectories are followed during the heating and during the cooling. Comparing at the same temperature close to room temperature, a test piece which has finished the heating and cooling contracted by 25 ppm (hereinafter, referred to as “amount due to seasoning effect”) from its initial size. Before this measurement, the test piece had been subjected to a heat treatment at 98° C. for 48 hours. Subsequently, this measurement was again carried out after 570 hours had elapsed at room temperature from the measurement. As a result, among the carbon atoms formed in solid solution, a certain proportion can be thought to have moved to a stable position at room temperature (hereinafter, referred to as “room temperature stable position”. Likely to be a tetrahedral interstitial site in a fcc lattice). If the heating is carried out to 205° C., spontaneous magnetization almost completely disappears, so that at this temperature carbon atoms can be thought to move to another position (hereinafter, referred to as “high-temperature side stable position”. Likely to be an octahedral interstitial site in a fcc lattice). Further, it is considered that during cooling the diffusion of carbons does not keep up, whereby the material returns to room temperature with most of the high-temperature side stable positions (a position where the carbon atoms are stable at 205° C., not room temperature) as is.

Thus, it can be assumed that the above-described 25 ppm contraction occurred as a result of the change from the volume (expanded state) of the test piece in a state where a certain proportion of the carbon atoms were moved to a stable position at room temperature to a volume (contracted state) in a state where carbon atoms were moved to an unstable position at room temperature due to the heating and cooling of the first cycle.

The temperature of 98° C. appears to have been determined for the purposes of increasing the rate of interstitial diffusion of the carbon atoms and finishing the treatment at a time which is practical for an industrial manufacturing process. It was probably thought that at this temperature, since the decrease in spontaneous magnetization is still small, most of the carbon atoms would be at the same position, which is a stable position at room temperature. For this point, as described below, the diffusion of carbons from a room temperature stable position to a high-temperature side stable position appears to start at 80° C. or more in the present investigation. A temperature of 80° C. is thus thought more suitable than the above-described artificial seasoning temperature of 98° C.

Although the optimum treatment for temporal deformation is to leave a member at a use temperature for several tens of years and then fit the member into a product, this is not practical. However, in the present investigation, there are also materials for which measurement of the third cycle was performed after leaving at room temperature for 1349 hours, without carrying out the 98° C. heat treatment between the second cycle and the third cycle. In this case, the amount due to the seasoning effect was 19 ppm. In other words, during this time temporal deformation of a 19 ppm expansion occurred. Since 1349 hours is not a practical time period for an industrial treatment, the below-described temperature range is better for the stabilizing treatment. When the second cycle measurement was carried out immediately after the first cycle was finished, the size of the test piece was not changed even after undergoing heating and cooling. This can be assumed to be because at the first cycle heating temperature (210° C.), all of the carbon atoms had diffused to the high-temperature side stable position.

In the second cycle, the reason why the curves do not exactly match during the heating and the cooling is that the temperature change of the test piece is slower than the temperature change of the control thermocouple. As a result, the size difference is larger in the range where the coefficient of thermal expansion is large (high-temperature side). To confirm this, in a graph of the measurement results for a not-illustrated low-carbon material (C: 0.002 wt %), both curves matched when displayed with an offset of +4° C. on the heating curve and −4° C. on the cooling curve. In the measurement of this diagram, a stabilizing treatment at 98° C. for 48 hours was subsequently carried out, then the material was left at room temperature for 816 hours until the measurement and the third cycle measurement was carried out.

Similar to the measurement results of the first cycle, in the third cycle a fairly large amount due to the seasoning effect was again manifested. Compared with the 25 ppm in the first cycle, this was 30 ppm in the third cycle. This is thought to be due to the long standing time at room temperature.

The stabilizing temperature must be higher than room temperature and be at least less than the Curie temperature. However, since the actual Curie temperature should basically be at the high-temperature stable position, a temperature of 150° C. or less, which is where the heating curve and the cooling curve intersect, is realistic. In the present measurement, since it took about 2 hours to heat from room temperature to 150° C., it can be said that there is an artificial seasoning effect even with a treatment of 150° C. for 2 hours or less.

Further, while it can be understood from the heating curve, the temperature at which the heating curve and the cooling curve (straight portion of 100° C. or less) are no longer parallel is 80° C. At temperatures higher than that, these curves are not parallel. From this result, it can be estimated that the carbon atoms present in room temperature stable positions at 80° C. or more start to diffuse to high-temperature stable positions. Therefore, to make most of the carbon atoms be in room temperature stable positions, it is more desirable for the temperature to be above room temperature to 80° C. or less.

The fourth cycle followed the same trajectory as the second cycle. Since measurement is carried out immediately after the third cycle, and the stabilizing treatment is not carried out, this is a projected result.

From the above, a member which has finished the second to fourth cycles will have undergone the same treatment as gradually cooling from a temperature of 205 to 210° C. However, the thus-treated alloy should not be fitted into a product. This is because, for a high-carbon material, temporal deformation which is larger only for 25 to 30 ppm than a member which has undergone a stabilizing treatment will occur in the future.

Further, the fact that the high-carbon materials are subject to temporal deformation of expanding by 18 ppm calculated on an annual basis was described above. However, this measurement of temporal deformation was carried out on a material after it had already undergone a stabilizing treatment. Thus, this does not contradict the fact that the same amount due to seasoning effect is produced even for a material left at room temperature for just 1359 hours (about 2 months).

While the dimensional change of a low-carbon material test piece was measured in the same manner as for the high-carbon materials, regardless of whether the test piece was left at room temperature before measurement or not, the sample length returned to the initial size even after the heating and cooling. Specifically, the contraction which was seen for the high-carbon material could not be confirmed. Therefore, as mentioned above, the reason why the slope becomes smaller at 80° C. in the heating curve of the high-carbon material and the size is contracted at the point when cooling was carried out to room temperature can be assumed to be due to the diffusion of carbon atoms.

While the present invention has been described with reference to exemplary embodiments, it is to be understood that the invention is not limited to the disclosed exemplary embodiments. The scope of the following claims is to be accorded the broadest interpretation so as to encompass all such modifications and equivalent structures and functions.

This application claims the benefit of Japanese Patent Application Nos. 2008-117352, filed Apr. 28, 2008 and 2009-097226, filed Apr. 13, 2009, which are hereby incorporated by reference in their entirety. 

1. An alloy comprising iron, nickel, and cobalt which are basic components of a Super Invar alloy, characterized in that an amount of a fraction which has not been carbidized in carbon contained in the alloy is 0.010 wt % or less.
 2. The alloy according to claim 1, characterized by further comprising carbide forming elements, wherein carbides formed from at least a part of the carbide forming elements and carbon are dispersed in a base phase.
 3. The alloy according to claim 2, characterized in that a total carbon content in the alloy is 0.010 wt % or more.
 4. The alloy according to claim 3, characterized in that the carbide forming elements are contained in a higher amount than the total carbon content in terms of atomic percent.
 5. The alloy according to claim 2, characterized in that a content of the carbide forming elements is 0.05 wt % or more and 0.50 wt % or less.
 6. The alloy according to claim 2, characterized in that an amount of carbide forming elements which have not been carbidized is 0.50 wt % or less.
 7. The alloy according to claim 2, characterized in that a compound phase of the nickel and the carbide forming elements is formed.
 8. The alloy according to claim 2, characterized in that the carbide forming elements are titanium or niobium.
 9. A lens holding member comprising the alloy according to claim
 1. 10. An optical apparatus comprising the lens holding member according to claim
 9. 11. A method for producing an alloy comprising iron, nickel, and cobalt which are basic components of a Super Invar alloy, characterized by comprising: adding carbide forming elements to the basic components and melt casting the resultant mixture; hot forging at a predetermined temperature; and precipitating carbides formed from carbons contained in the alloy and the carbide forming elements in a base phase by performing a first heat treatment at a first temperature which is lower than the predetermined temperature.
 12. The method for producing an alloy according to claim 11, characterized in that the first temperature is 825° C. or more and 950° C. or less.
 13. The method for producing an alloy according to claim 11, characterized in that a content of the added carbide forming elements is 0.05 wt % or more and 0.50 wt % or less.
 14. The method for producing an alloy according to claim 11, characterized by further comprising forming a compound phase of the nickel and the carbide forming elements by performing a second heat treatment at a second temperature which is lower than the first temperature.
 15. The method for producing an alloy according to claim 11, characterized by performing a third heat treatment at a third temperature which is lower than a Curie temperature of the alloy.
 16. The method for producing an alloy according to claim 15, characterized in that the third temperature is 80° C. or less. 